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Technical Report No. 8

*Electrically Conducting Polymers

by

W. D. Gill, T. C. Clarke, and G. B. Street

Prepared for Publication

in

Electrical Properties of Polymers

(A. M. Herman, ed.)

IBM Research Laboratory5600 Cottle Rd.

San Jose, CA 95193

April 7, 1983

*O Reproduction in whole or in part is permitted forany purpose of the United States Government

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I Research ReportELECTRICALLY CONDUCTING POLYMERS

W.D. GilT. C. ClarkeG. B. Street

[EM Research LaboratorySan Jose, Cafifornia 95193

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ELECTRCALLY CONDUCTING POLYMERS

W. D.GillT. C. ClarkeG. B. Street

IBM Research LaboratorySan Jose, California 95193

ABSTRACr

* This paper reviews the present state of the field of conducting polymers.

1

L INTRODUCTION

In parallel with developments in the field of quasi one-dimensional organic metals

[1-3], intense interest in other possibly low-dimensional conducting systems has led to

investigation of a number of polymer systems with high conductivity potential. The

obvious chemical and structural anisotropy of polymers immediately suggested that a

conducting polymer might have the highly anisotropic electrical properties seen in quasi

one-dimensional conducting crystals such as KCP [2] and TTF-TCNQ [1,3].

* Polycrystalline compactions of polythiazyl (SN) x were known to be highly conducting

from the work of Chapman et al. [4], who had concluded that the material was a

semiconductor. In 1973 Walatka et al. [5] reported results of electrical measurements on

crystals of (SN),. The metallic conductivity along the polymer chain axis observed by

these authors stimulated extensive research into the properties of the material. With the

discovery of superconductivity in (SN), crystals by Greene, Street and Suter [6] the

investigation of (SN) x and the search for other highly conducting polymers rapidly

developed.

Attempts to synthesize chemical analogs of (SN), were disappointing, but

modification of the electronic properties of (SN) x by reaction with oxidants led to the

interesting halogen modifications of (SN) x by Bernard et al. [7] and independently by

Street et al. [8,9] and Aktar et al. [10,11]. The brominated (SN) 1 modification

(SNBr 0 .4 )x with conductivity of 2-3 x 104 ohm" -cm-1 remains the highest conductivity

polymer measured to date.

The next advance in this field was the discovery of organic polymer systems with

metallic properties. For many years the properties of long chain polyenes had been

2

theoretically investigated as potential semiconductors (12]. However, the longest polyene

chains were less than 20 units long. As far back as 1958 polymerization of acetylene to

essentially infinite polyene chains had been successfully carried out in the presence of a

Ziegler catalyst (13]. The product of these early reactions was a fine powder of

polyacetylene, whose powder compaction conductivity was found to increase substantially

on exposure to a wide range of oxidizing molecules [14]. In 1974, Ito et al. (15]

discovered that by exposing high concentrations of the Ziegler catalyst to a relatively high

pressure of acetylene gas, polyacetylene films were formed. Then in collaboration with

MacDiarmid and Heeger at the University of Pennsylvania doping experiments on

polyacetylene films with various oxidants resulted in conductivities as high as

1000 ohm' 1-cm-1 [16]. This discovery has stimulated extensive investigation of

polyacetylene and other organic polymers in many laboratories with the result that a

number of organic polymers with metallic-like conductivities have now been discovered.

Our purpose in this chapter is to describe some of the physics and chemistry of thise

highly conducting polymer systems ranging from the inorganic (SN) x to organic materials

such as (CHY1 , polypyrrole and the polyphenylenes (171. It would be very difficult in a

* single chapter to adequately review the already extensive literature on these materials.

Instead we hope to-present a coherent overview of the current status of highly conductive

polymers and-to-referthe-ineested-rmader to several review articles for a more extensives.:4- survey of the literature.

o'p' '

3

EU. POLYTHIAZYL (SN) x

A. Synthesis of (SN) x

1. (SN),, CRYSTAL GROWTH, FILM GROWTH AND STABILITY

In his paper in 1910, Burt (18] true to his middlename, Playfair, acknowledged that

(SN) x was first prepared by Davis. Burt prepared both films and crystals of (SN) x by the

thermal pyrolysis of S4 N4 in the course of trying to remove excess sulfur from the S4 N4

by passing its vapors over silver wool. The silver wool acts as a catalyst [19] in the

conversion of S4 N4 to S2 N2 (20]. As the crystal structure of S2 N2 reveals [21] the

square planar S2 N2 molecules are ideally arranged for a solid state polymerization to

(SN),. The mechanism of the solid state reaction has been studied by Baughman et al.

[22]. Burt's paper provides the basis of the techniques used today to prepare the (SN) x

crystals which have been used in the investigation of its electrical properties. All of these

techniques [23-25] emphasize the importance of the purity of the intermediate S2 N2

* crystals and careful temperature control during their slow solid state polymerization.

' Labes et al. [26,27] have also investigated crystal growth techniques involving

photoinitiation of polymerization of solution grown S2 N2 crystals.

Films of (SN) x have also been investigated extensively. They have been prepared

by two techniques, pyrolysis of S4 N4 and direct heating of (SN) x [28-30]. Films in which

the (SN) x chain direction is ordered can be prepared by deposition on mylar treated to

provide a series of parallel scratches on the surface [31]. The vapors above both (SN)1

and S4N4 have been studied in detail [32,33] and they contain species which would

contaminate the resulting films. This contamination can be minimized and the properties

of the films modified by careful control of the substrate temperature [34,35]. Smith er al.

(36] and Saalfeld et al. [37] concluded that the main component of the vapor from (SN) x

, . . _ - .. " .°.- *-.. " %' .o% . .- "

U. -"•% % ,% .', 'o' ." '.% -q','

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was an (SN) 4 bent chain species. Such a structure would have an electric dipole moment,

which is not consistent with the results of molecular beam electric deflection analysis [38].

Iqbal er al. [39] have interpreted IR spectra of (SN) x films deposited at low temperatures

as evidence for linear (SN) x oligomers.

Though both films and crystals of (SN), are stable indefinitely in inert atmospheres

they decompose in air [40]. This decomposition is observed most readily for thin films

and is also frequently first observed on those crystal faces where the chains terminate.

The stability of the films and crystals is negatively affected by the presence of

unpolymerized S2 N2 as well as other impurities. Nuclear relaxation studies [41] have

shown that (SN) 1 adsorbs water quite rapidly, resulting in the presence of high hydrogen

concentration in the material. Seel et al. [42] have calculated the effects of hydrogen on

the density of states in (SN),. However, no experimental correlations exist between the

physical properties of (SN) x and its hydrogen content either from adsorbed water or other

sources.

B. Structure of (SN),

On the basis of its complete insolubility, Burt in 1910 characterized (SN)x as a

polymer [18]; however, it was not until 1956 that Goehring and Voigt first investigated its

structure and suggested a zig-zag geometry for the chain [43]. Doullard [44] showed that

this geometry was incorrect and advanced what has proved to be essentially the correct

chain structure. A preliminary crystal structure obtained from electron diffraction [45]

using this chain geometry was later refined by X-ray diffraction [21] and confirmed by

neutron diffraction techniques [46]. This structure is shown in Figure 1. The unit cell

contains two almost flat chains, which are centrosymmetrically related but translationaly

. . .. *

S.w 5

inequivalent. These inequivalent chains alternate along the c-axis whereas equivalent

chains are adjacent along the a-axis. The chains all lie in the 102 plane where equivalent

chains alternate. This 102 plane is an easy cleavage plane. The shortest interchain bonds

are tabulated in Table 1. Comparison of these bond lengths with the corresponding

van der Waals diameters suggests that the interaction between the chains are weak

relative to the bonding within the chains.

Studies of the vibrational spectra of (SN) x using both Raman and IR techniques

[47-501 have also given useful insight into the relative strengths of the intrachain and

interchain bonding. In general these data confirm the conclusions from other

measurements that (SN), has a high degree of structural anisotropy. The intrachain force

-" constants are significantly larger than the interchain force constants.

C. Electronic Properties

1. BAND STRUCTURE

The intrinsic electronic properties of any crystalline solid, i.e., whether it is a metal,

semimetal, semiconductor or insulator, are determined by its electronic band structure.

Once the crystal structure has been completely determined, this band structure can in

principle be calculated exactly. However, due to the large number of electrons in the unit

cells of polymeric conductors (44 electrons in the case of (SN) x) various approximations

are necessary. A large number of band structure calculations for (SN), have been made

ranging from relatively simple one-dimensional extended HUckle tight-binding methods to

a complete three-dimensional orthogonalized plane wave calculation [51].

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The (SN), structure contains two chains per unit cell. each with two SN pairs

related by a screw axis symmetry. Simple one-dimensional band calculations for this

structure [521 predict a half-filled conduction band with a degeneracy at the Brillouin zone

boundary. This degeneracy should be split by a Peierl's distortion resulting in an

insulating state for the system. Experimentally it was known that (SN) x was metallic and

did not undergo a Peierl's transition even at very low temperature [5]. It was therefore

necessary to include interchain interactions which could stabilize the structure against such

a Peierl's distortion. A variety of three-dimensional calculations have been made which

show that quite extensive interchain coupling in (SN) x stabilizes the metallic state; thus

(SN) 1 must be regarded as a highly anisotropic conductor instead of a quasi

one-dimensional metaL As we will discuss below these calculations verified the

."".-. conclusions on dimensionality already experimentally established from optical anisotropy

[53] and electron energy loss measurements [54].

The three-dimensional OPW calculation of Rudge and Grant [55] gives the most

complete picture of the band structure (Figure 2) and Fermi surface of (SN),. The Fermi

level crosses the overlapping bands in the chain direction (1 to Z) and in the

* perpendicular direction (Z to E), resulting in electron pockets at Z and hole pockets at E.

.. Thus, the calculations show that (SN) x is a semimetal, but perhaps more importantly they

also demonstrate that although (SN) x is strongly anisotropic, it should not be regarded as

a quasi one-dimensional conductor. The expected degree of anisotropy has been estimated

- from the OPW band structure by Grant et al. [56], who derived the plasma tensor

anisotropy by suitable integration over the Fermi surface. The magnitude of the principal

axes for the plasm.: --Por ' :n the ratios 1:0.13:0.09 in the b:('02) 1:(T02) directions,

in reasonable agreement with experiments discussed below.I-So-

,/7

7

2. NORMAL CONDUCTIVITY AND SUPERCONDUCTIVITY

The transport measurements of Walatka et al. [5] in 1973 on (SN), crystals

stimulated the recent interest in this material. Conductivity and thermoelectric power

measurements strongly suggested that (SN) x was metallic to 4.2K in the direction parallel

to the (SN) x chains in the crystal structure. Strong anisotropy was observed for

conductivity parallel and perpendicular to the fiber axis of the crystals, which coincides

with the (SN), chain axis. This anisotropy led to early speculation that (SN)1 was a quasi

one-dimensional conductor. However, the persistence of metallic conductivity to very low

temperatures, where a Peierl's transition to the insulating state would be expected for a

one-dimensional system, suggested that the fibrous nature of (SN) 1 crystals might be

dominating the observed conductivity anisotropy.

In 1975, Greene et al. [6] discovered superconductivity in (SN) x with Tc=0.3K.

This observation, which dramatically verified the metallic properties of (SN) x, was also the

first observation of superconductivity in a polymer and the first observation of

superconductivity at ambient pressure in a material containing no metallic elements. The

superconductivity experiments suggest that (SN) x is a bulk type II superconductor but that

crystal perfection strongly affects the superconductivity.

The effects of improved crystal perfection on the conductivity properties is shown in

Figure 3 [57]. The maximum conductivity increased from 1000 Q'1cm-1 to about

4000 Q'1cm'1 while the resistivity ratio PRT/P4K increased from about 3 to 250. The

superconductivity transition temperature Tc also increased from 0.26K to 0.35K. The

quadratic dependence of resistivity on temperature observed with the best crystals

suggested to Chiang et al. [58] that the usual electron-phonon scattering processes are not

. ., . . .. -.- -. ,* •- . . . . , - 4. .*** .... .. - *,,; , .,,,,. .?.," ,

80

dominant in (SN) x but that electron-electron Umklapp scattering between electron and

hole pockets of the Fermi surface is dominant.

The unusual pressure dependence of normal conductivity should be mentioned here

since it provides additional support for the electron-electron scattering process and also

gives some insight into the transport and optical properties of brominated (SN) x described

in a later section. Under hydrostatic pressure the increase in normal conductivity is more

than an order of magnitude greater than expected from lattice stiffening effects on

electron-phonon scattering processes [59]. However, electron-electron scattering can be

expected to be extremely pressure sensitive through its dependence on details of the Fermi

surface. Consideration of the effects of hydrostatic pressure on the band structure

" indicate that the electron pocket may be shifted from the Brillouin zone center thus

reducing the electron-hole scattering probability with a resultant increase in conductivity

[651.

Early attempts to observe a Meissner effect in (SN), crystals were unsuccessful (611

leaving some doubt as to whether the observed superconductivity was a bulk effect.

However, a broad hump in the specific heat indicative of the expected BCS anomaly was

observed (62]. More recently the Meissner effect has been observed [63,64] verifying the

model of (SN) x as a bulk, weakly coupled filamentary Type II superconductor. A number

" of interesting problems related to superconductivity in (SN) x still need to be resolved.

These include the pressure dependence of Tc, which increases strongly with pressure to

about 8 Kbar where a sudden decrease in Tc apparently signals a change in phase [65,661;

the absence of superconductivity in (SN) x films [67-69]; and the absence of al.u

-2" superconducting gap in tunneling experiments [70]. Full understanding of the temperature

......................................... ... ..........................

9

dependence of the critical magnetic field and its strong anisotropy [71] also provides

interesting challenges. The explanation of the very broad superconducting transition width

as fluctuations of one-dimensional superconductivity [72] is intriguing in view of the

results observed in (SNBr 0.4)1 described below.

3. OPTICAL PROPERTIES

The optical properties of (SN) 1 crystals have been measured by several groups

[53,73] and have played an important role in understanding the intrinsic properties.

Optical reflectivity with the E-vector of the incident light parallel to the chain axis

(b-axis) shows a steep, metallic reflectance edge which can be analyzed using the Drude

model,yielding a plasma frequency hpp-6.9eV and a scattering time -=2.6 x 10"15 sec

[53]. The optical conductivity derived from these measurements,

aopt' 2 /4- 2 .5 x 104 Q' 1 cm' 1 , is about an order of magnitude greater than the measured

DC conductivity. With the E-vector perpendicular to the b-axis a much weaker structure

is observed in reflectivity and is interpreted to be a strongly damped plasma edge at

hpp-2.4eV with T"=3 x 10- 16 sec. The observation of this perpendicular reflectivity edge

was the first experimental evidence that (SN) , although strongly anisotropic, should not

. be regarded as a quasi one-dimensional metal. This conclusion is also supported by

electron energy loss experiments (54] and by magnetoresistance anisotropy [74] and is in

good agreement with the three-dimensional OPW band structure calculations [55,56].

4. OTHER PHYSICAL PROPERTIES

A broad range of experimental techniques have been applied to (SN)x to measure

both its lattice and electronic properties. These experiments include specific heat [751,

inelastic neutron scattering [76], IR reflectance, Raman scattering [47,48], X-ray (77] and

10

ultraviolet photoemission [78], magnetoresistance [74,79,80], Hall effect [801 and junction

properties (81,82]. Space permits only a very brief review of these experiments which,

however, demonstrate some very interesting and important properties of this material.

From the electronic contribution to the specific heat, Harper et al. [751 obtain a

value for the Fermi level density of states of 0.14 states/eV-spin-molecule in excellent

agreement with band structure predictions. The temperature dependence of the lattice

contribution to the specific heat is characteristic of a solid with quasi one-dimensional

binding forces. From IR, Raman and neutron scattering results, Stolz et al. (47] estimate

* that the interchain force constants are about ten times smaller than the intrachain force

constants.

The X-ray photoemission (77] and uv-photoemissior f-78] results give valence band

densities of states which again agree very well with band structure predictions (see

Figure 2). A value of about 0.5 electrons transferred from S to N is obtained from these

experiments consistent with theoretical predictions.

Several measurements of the galvanomagnetic properties of (SN) 1 crystals have been

reported. Beyer et al. [74] measured the magnetoresistance anisotropy at 4.2K and

compared it with the anisotropy derived from the calculated plasma tensor of Grant et al.

[6]. A negative magnetoresistance dominated the measurements at low magnetic fields.

At higher fields a quadratic positive magnetoresistance was observed with a parallel to

O transverse mobility anisotropy ratio of 3.1, in reasonable agreement with calculated value

of 5.7. The b-axis mobilities are about 600 cm' 2 /V-sec and 400 cm 2/V-sec for holes

and electrons, respectively. A scattering time of 1.5 x 10,13 sec was obtained

:..:. ...,,.,.....,..,. .. ,. -....-. .

" " 11

corresponding to a mean free path of -700A along the b-axis. The Hall effect could not

be observed in these experiments. Kahlert and Seeger [80] have reported

magnetoresistance on samples with exceptionally high ratios of 04K/O300K. No negative

component of magnetoresistance was observed until after severe bending of the crystals,

suggesting that this anomalous effect is connected with defects. Kaneto et al. [79]

attribute the negative magnetoresistance to spin-dependent scattering at chain breaks or

defects with localized spins. Kahlert and Seeger [80] have observed the Hall effect for an

(SN), crystal, from which a value of the effective carrier density neff- 3 .2 x 1021 cm 3

was obtained in excellent agreement with band structure results.

Junction properties measured by Scranton et al. [81] have shown the interesting

result that (SN) 1 is more electronegative than gold, producing large Schottky barriers on

contact with n-type semiconductors. A practical utilization of this property in Schottky

.arrier solar cells has been investigated by Cohen and Harris [82].

5. (SN) x FILMS

The previous discussion has been concerned with the properties of (SN) 1 crystals.

*O" However, thin films of (SN), have interesting optical properties and a wide range of

electrical properties.

Although the polycrystailine films are usually disordered, films with highly

- anisotropic optical properties were obtained by Bright et al. (3 1 from depositions on

stretched films and especially on mylar films which had been lightly abraded or rubbed in

one direction. Films deposited on these substrates have almost complete alignment of the

(SN) x crystallites with the polymer chain axes (b-axis) in the direction of the stretching or

.O

12

abrasion. This crystallite alignment results in strongly anisotropic optical and transport

properties. The optical anisotropy has interesting potential application as a polarizer in

the visible and near infrared.

The electrical properties of (SN) x films have been studied extensively [67,68,29,35].

The conductivity is dominated by interparticle barrier effects resulting in an activated-like

temperature dependence instead of metallic-like behavior. The effects of substrate

temperature on conductivity, thermoelectric power, Hall effect and magnetoresistance have

been investigated by Beyer et al. [68,35]. Thermoelectric power and conductivity results

0. shown in Figure 4 indicate that the metallic transport properties of (SN) x crystals are

approached as the substrate temperature is increased. However, superconductivity was

not observed even in the most metallic of these films. A model for the film transport

results suggests the existence of a gap in the density of states correlated with (SN), chain

length (35]. Such a gap might also explain the absence of superconductivity in (SN) x

films [68].

II. CHEMICAL MODIFICATIONS OF (SN) x

*i A. General Survey

The discovery of superconductivity in crystals of (SN) x motivated intensive efforts

in several laboratories to synthesize analogous compounds. Much of this work was

focussed on efforts to produce (SeN), from the known compound Se4 N4 by routes similar

to those described for conversion of S4 N4 to (SN) x [57]. These efforts have not been

successful. Similarly unsuccessful attempts have also been made to produce (SCN) x [83].

Oligomeric analogs of (SN) x such as 1,9-diaryl-pentasulfurtetranitride [84] have been

prepared but they showed no evidence of conductivity despite the fact that there was an

|y............- .. . .

° . . . . . . . .A . * -, . , .* ~ . , . o •,"" " . '" - " . " ,' " -_ " * " . , * *" ,. ." ," . . .-' "J - *. . . . . . . . . .i -t ": " ¢ m , -

13

• .increasing red shift of the uv-visible absorption with increasing chain length indicating

extensive electron delocalization along the chain. Attempts to make even longer chain

* oligomers were not successful.

B. Halogen Modifications

1. PREPARATION

The (SN), chain and crystal structures were shown above in Figure 1. The unit cell

*' of (SN), contains two, almost fiat, translationally inequivalent, centrosymmetrically

related (SN), chains. These two inequivalent chains alternate along the c-axis, whereas

equivalent chains are adjacent along the a-axis. The chains all lie in the 102 plane where

equivalent chains alternate. The ready cleavage which takes place at these 102 planes

suggest the possibility of intercalating various molecules between them. In 1976 Bernard

et al. [7] reported the formation of an intercalation compound of (SN)1 and bromine.

Independently Street and Gill [8,9] surveyed a large number of potential intercalates and

found that the halogens Br 2 12, ICI and IBr all increased the conductivity of (SN),. The

bromine derivatives have been more extensively studied because they show the highest

conductivity.

The simplest preparative technique for brominating (SN), is exposure to dry

bromine vapor. On exposure (SN) x crystals change color from gold to blue-black and

,' expand in directions perpendicular to the chain axis. Exposure to the vapor pressure of

bromine at room temperature leads to a composition (SNBr 0 .5 )x which on pumping in

vacuum (10- 5 Torr) at room temperature for one hour gives a final composition of

(SNBr 0.4) x . At this composition the crystal has expanded by approximately 50% in

2* volume perpendicular to the chain axis. The flotation density increases from 2.32 g/cm3

tI

L5 14

K.-: to 2.65 g/cm3 . Virtually any composition of lower bromine content can be obtained by

heating a sample of (SNBr 0.4 ) x in vacuum for various periods of time. However, the

(SNBr0 .4 )x composition appears to have the best electrical properties and has been most

extensively studied. It is important to brominate with bromine vapor because (SN) 2 reacts

with liquid bromine yielding significant amounts of a side product S4N3 Br 3 (851.

A different preparative route involves direct bromination of the (SN) 1 precursor

molecule S4 N4 [86,87]. Exposure to bromine vapor causes S4N4 crystals to turn black

and expand rapidly destroying their external habit. After removal of free bromine the

* . remaining, conductive black powder has the composition (SNBr 0 4 )x- IR [88], Raman

[89], thermal analysis [88] and mass spectrometry [90] all indicate that this material is the

same as that obtained by bromination of (SN) x crystals; however, it is even more

structurally disordered. The mechanism of this ring opening reaction is unknown. ICI and

"Br vapor also react readily with S4 N4 to form conducting solids [86,87,91].

Considerable effort has gone into studies of the molecular nature of the intercalated

halogens in (SN), derivatives. In graphite, bromine is known to be present in the form of

0. weakly ionized Br 2 (92]. However, in (SN) x Raman spectroscopy indicates a much more

complex situation (89,50]. Two strong Raman fundamentals are observed at 150 cm "I

and 230 cm 1 . The first band has been assigned without ambiguity to the symmetric10

stretch of the tribromide ion, Br-. The band at 230 cm' t can be assigned to the

. asymmetric stretch of the same Br3 ion or to the stretching frequency of Br 2 downshifted

* from its 325 cm "1 frequency in the gas phase by its strong association with the (SN) x

lattice. Macklin et al. [93] examined the IR absorption of brominated (SN) x and found

no evidence for the 230 cm"1 peak which would be expected to be IR active for the

4! 15

asymmetric stretch of a linear Br" ion. On this basis it is concluded that both Br 2 and

Br" molecular species are present in brominated (SN),. From the polarization

dependence of the Raman peaks the Br 2 and Br" appear to be oriented with their axes

parallel to the (SN), chains. Extended X-ray Absorption Fine Structure (EXAFS) studies

[94] are also consistent with this orientation and indicate that similar concentrations of

Br 2 and Br3 species are present in the (SNBr 0.4 )x crystals. Magnetic susceptibility

measurements show the absence of paramagnetic species such as Bri [95].

Mass spectrometric studies of the species emanating from the surface of heated

;- (SNBr 0 .4 )x [90,96] show an initial high concentration of Br 2 which rapidly decreases and

a lower concentration of SNBr which becomes the dominant volatized species after the

first few hours. These results can be rationalized by assuming the bromine to be in the

two molecular forms Br 2 and Brj.

2. STRUCTURE OF (SNBr 0 .4 )x

The incorporation of bromine into the (SN) 1 lattice increases the disorder and

* twinning so that a complete single crystal X-ray structure determination is not possible

. [9,54]. Data from which the mode of bromine incorporation is inferred have been

obtained from electron diffraction [9], X-ray powder diffraction [7,9,50], EXAFS [941,O

Raman [89] and IR spectroscopy [93]. In their original experiments, Bernard et al. [7]

concluded from X-ray powder data that only the repeat unit in the chain direction remains

unchanged on bromination. This result has been confirmed from X-ray precession

photographs [8,9]. Iqbal et al. [50] showed that as the bromine content increased, d(100)

• 4 . . • . . . .- -- . . ." 4 , 5 . " - * S % . % ' % '' - . % % % S" t-'h S% %S0

~~~~~~~~~~. .. . . .. .. . . .+4 .. .= - ,- . - * . . *+ . - -. , . .. • -.z 7 z) 7 . 77 ZL 7. 7.°.o.. .+° - ••-. .",-•• °

"-.

16

0O

decreased and d(002) increased to an essentially common value of 3.62A. The interplanar

separation d(102) was found to decrease.

Electron diffraction studies show patterns typical of (SN) 1 but with considerable

streaking perpendicular to b* indicating increased disorder and (100) twinning (Figure 5)0 0

[9]. Bromination results in a decrease in the average fiber diameter from -70A to -30A.

Superimposed on the (SN)x-like diffraction pattern are diffuse fines at (0,1/2k,0)

corresponding to a one-dimensional commensurate superlattice oriented along the chain

axis with a repeat unit twice that of (SN) x . In addition, diffuse lines appear near 2b/3,

2b/5 and 2b/7 (971. Diffuse X-ray scattering measurements by Comes (98] confirm

these electron diffraction observations. On cooling below - 150K the 2b superlattice lines

disappear, but they reappear on warming (97]. This reversible transition is quite sharp,

occurring over about a five degree temperature interval. However, the diffuse lines

remain at low temperature but irreversibly weaken in intensity on heating the samples

above 100C. This behavior suggests that these diffuse lines are associated with the

adsorbed Br 2 species, which is known to evolve at these temperatures from the mass

spectrometric results. The 2b superlattice is believed to be associated with ordering of the

linear Br" ions along the (SN) x chains. The disappearance of this superlattice structure at

low temperature is not understood. On the basis of Raman studies of iodinated (CH) x,

Fubino et al. [98] have suggested that the iodine is in the form of 13, 15 and 12.

Attempts to understand the structure of ICI/(SN), complexes have met with much more

limited success, but both Raman- and electron-diffraction results (97] have been

interpreted on the basis of un-ionized ICI adsorbed on the surface of the (SN) 1 .

"0

= 4

17

3. ELECTRONIC PROPERTIES

Bromination of (SN) x results in marked changes in optical properties, a dramatic

increase of the normal conductivity, and improved superconductivity characteristics. It is

-. particularly interesting that these improved properties can be attributed to subtle changes

*in (SN), characteristics due to electronic effects of the intercalate. The electronic

.- properties of brominated (SN)1 suggest that this material can, to first order, be regarded

- as (SN), whose transport properties have been optimized by the bromine interaction.

The conductivity of (SN) x films and crystals exposed to bromine to form (SNBr 0 .4 )x

is increased by an order of magnitude so that a 2 to 4x 104 U-lcm-1 at 300K [8-11].

On cooling, the temperature dependence is weaker than the T2 dependence of (SN)x .

Thermoelectric power measurements (9] show a change in sign on bromination from

negative to small positive values consistent with bromine acting as an acceptor. The

conductivity perpendicular to the (SNBr 0 .4 )x fibers, unlike that of (SN) x , is metallic to

3K, below which' it is temperature independent [9]. This behavior suggests that

(SNBr 0 .4 ) x is more three dimensional, i.e., the fibers are better coupled than (SN) x itself.

In contrast to (SN)1 , where a increases dramatically under pressure [59], Gill et al. [16]

report only a 10 to 20 percent conductivity increase under 10 Kbar.

The optical properties show that a strong red-shift of the plasma edge occurs on

bromination (see Figure 6) [9,10,11]. However, Drude-Lorenz analysis indicate little or

no change in plasma frequency or optical scattering time. The shifted plasma edge, which

accounts for the observed color change, is accounted for by an increased background

dielectric constant on bromination.

18

The effects of bromine incorporation on the superconducting properties are

Particularly interesting since the transition temperature TC changes very little [91, while the

width of the transition, the pressure dependence of Tc (99], the Meissner effect (63] and

V. the critical field are strongly affected (105]. Initial measurements of TC on (SNBrO.4 )1

* crystals showed negligible change from (SN), values [9]1. More recent results show an

increase of about 20 percent in TC on bromination (100]. The transition is much sharper

in (SNBrO. 4)x than in (SN)1 as can be seen in Figure 7. The broad transition in (SN) 1

'O.

has been partly ascribed to disorder (511; however, since (SNBre. 4 ). has considerably

more disorder and a sharper transition, the role of the disorder appears not to be dominant

in either material. The width of the transition in (SN) is more probably due to

superconducting fluctuations, whose absence in (SNBrO.4 )1 is evidence for increased

three-dimensionality in this material. The pressure dependence of Tc is also qualitatively

-. different from (SN) 1 (Figure 8), decreasing monotonically to about 150 mK under

10 Kbar. Dee et al. [631 have observed a complete Meissner effect in (SNBr.4) x

crystals. A much sharper transition than in (SN) 1 was also observed indicative of more

tightly coupled fibers. Critical field studies by Kwak et al [1001 show that in contrast to

(SN)o, H ) is well-behaved and can be treated by anisotropic three-dimensional

Ginzburg-Landau theory. These data are consistent with increased interfiber coupling in

(SNhBron4 )xr

Magnetic susceptibility measurements on brominated (SN) show a small Curie like

contribution below 30K [95]. An upper limit of 2x 10- molar for the concentration of

spin 1/2 species such as Br is obtained for (SNBr 4 )n. Above 30K susceptibility

increases linearly with temperature (95].0.

0-

19

Conductivity and magnetoresistance measurements on (SN), exposed to bromine,

iodine and ICI have been reported by Philipp and Seeger [101]. Only bromine

significantly improved the conductivity in these experiments. Halogen treatment enhances

a negative magnetoresistance component which is attributed to increased (SN)x chain

interruptions.

4. MODEL FOR (SNBr 0.4 )x

In the preceding two sections we have seen that although dramatic changes in the

appearance, structure and electronic properties occur on bromination of (SN),, closer

examination shows that the basic features (both structural and electronic) remain those of

(SN)x. From electron microscopy we see an almost unchanged (SN), diffraction pattern

with increased disorder and a superimposed superlattice structure. In the electronic

properties we see an unchanged plasma frequency and an almost unchanged

superconducting transition temperature. Clearly the model of (SNBr 0 .4 )x accounting for

these observations must retain many of the features of pristine (SN),. In the present

section we will outline the model for (SNBr 0 .4 ) which emerges from the wide range of

experimental observations.

From spectroscopic data we know that the bromine intercalate is incorporated as Br 2

and Bri, and that both species are aligned parallel to the (SN), chain axis. The decrease

-_ in d(102) indicates that bromine is not intercalating between (102) planes, but is enteringI.

substitutionally to occupy chain sites. Observation of the "2b" superlattice implies

periodic ordering of the bromine only in the direction parallel to the (SN) x chains.

Geometrical considerations suggest that Br3 ions are responsible for the superlattice

structure. Because bromine bonds to sulfur more readily than to nitrogen, the bromine

* . ,. y

20

molecular species can be expected to align along the chain axis so as to maximize

sulfur-bromine interactions and minimize nitrogen-bromine interaction.

Because of the fibrous morphology of (SN) 1 bromine will also be incorporated

between fibers. As the fiber diameter decreases, the amount of bromine attached to fiber

surfaces becomes significant relative to the amount intercalated within the fibers. For

(SNBr 0 .4)x with approximately 30A diameter fibers, all the bromine could be

accommodated as a monolayer on these fiber surfaces. Thus a structural model emerges

of 30A diameter (SN), fibers with a monolayer coating of aligned Brj/Br2 ordered along

the b-axis with a repeat of 2b due to preferential sulfur bonding conditions (24]. The

remaining Br 2 and Bri is thought to be incorporated within the fibers accounting for the

changed lattice parameters.

How does this model for the structure of brominated (SN), account for the dramatic

changes in the electronic properties? The large increase in conductivity and the

approximately constant plasma frequency suggest that the main effect of bromination is to

increase the dc scattering lifetime (9]. Increased free carrier density due to charge

0 transfer from free bromine cannot itself account for the tenfold increase in conductivity.

In (SN) x the resistivity has a T2 dependence suggesting that electron-hole scattering is the

dominant lifetime limiting process (58]. The Fermi surface geometry is consistent with an

efficient electron-hole scattering mechanism. Gill et al. [9] have suggested that the

increased conductivity of brominated (SN) x is due to suppression of electron-hole

scattering due to Fermi surface changes brought about by bromination. Assuming that

charge transfer from bromine removes approximately 0.1 electrons/SN unit from the

conduction bands, the Fermi level is lowered by about 1eV. Lowering EF strongly affects

S"A

. . . . . . . . . . . . ..-- . .

21

the shape of the Fermi surface, expanding the volume of the hole pockets and shrinking or

even eliminating the electron pockets. Such changes in the Fermi surface will certainly

drastically alter the special conditions favoring efficient Umklapp electron-hole scattering.

The assumed charge transfer of 0.1 electron/SN unit is consistent with the small plasma

frequency change observed optically. Elimination of electron-hole scattering through

acceptor doping with bromine increases the dc-lifetime and therefore the conductivity,

since now only the phonon scattering channel limits lifetime. The change from

electron-hole to electron-phonon scattering should be manifested in a change from T2 to a

linear temperature dependence for resistivity. In addition, one expects a much reduced

pressure dependence of conductivity in brominated (SN), since electron-phonon scattering

processes will not be as sensitive to the shape of the Fermi surface. In fact, the pressure

dependence of conductivity for brominated crystals is greatly reduced compared to (SN),

and can be adequately accounted for by lattice stiffening effects usually dominant in

metals. Thus this model invoking the Fermi surface changes due to charge transfer on

bromination and subsequent suppression of electron-hole scattering explains the major

features of the electronic properties of brominated (SN) x [60].

IV. POLYACETYLENE

A. Chemistry of Polyacetylene

Although polyacetylene, (CH)x, has long been of interest to theoreticians as the

o •ultimate member of the polyene family, the experimental study of this material was

impeded for some time by its intractable nature. Natta first reported a synthesis of

polyacetylene from acetylene using a Ziegler-Natta catalyst over 20 years ago, but the

polymer produced was a black, insoluble, air-sensistive powder with neither a melting

point nor a glass transition temperature [13]. Despite the materials limitations it was

X.1

22

recognized that this material might exhibit interesting transport properties. Thus, Hatano

carried out early electrical measurements which showed that (CH), is a wide band gap

semiconductor [102]. Even more interesting were the results of Berets and Smith who

showed that the conductivity of compressed pellets of polyacetylene could be

systematically varied from 10-9 to 10-2 ohm-l-cm " 1 by exposure to various Lewis acids

and bases [14]. These results aroused little interest, however, particularly in fight of the

dubious purity and uncertain structure of the polymer.

A major breakthrough in the synthesis and characterization of (CH) x occurred in

1971 when Shirakawa and Ikeda were able to form polyacetylene films by the

polymerization of acetylene on the surface of a highly concentrated Ziegler-Natta catalyst

solution (1031. These flexible films share the air sensitivity and insolubility of the earlier

Natta powders, but the physical continuity and convenience of handling of polyacetylene

in this new form provided the impetus for a more detailed investigation of the structure

and properties of this interesting material.

Films of polyacetylene exhibit a metallic luster and appear to the eye to be

continuous, but microscopic investigation of the morphology of these materials by

scanning electron microscopy reveals an open fibrillar structure [15]. The films consist of

interwoven fibers approximately 200A in diameter. Although the flotation density ofi. . polyacetylene is approximately 1.2g/cc, the bulk density is only 0.4g/cc [15,1041. Film

formation at high catalyst concentrations appears to involve a physical meshing of these

" growing fibers. Indeed, at lower concentrations of the same catalyst no film is formed,

apparently because the individual powder particles settle from solution before meshing of

the fibers can occur. By combined mechanical and thermal treatment Shirakawa has been

""* ** * * - - **

23

able to obtain partial alignment of the fibers in a polyacetylene film [1051. As expected

such stretch aligned films exhibit anisotropic electrical and optical properties [106,1071.

Shirakawa and co-workers examined a number of Ziegler-Natta catalysts, obtaining

optimum results with a 4:1 mixture of triethylaluminum and titanium tetra-n-butoxide

[15]. Subsequent investigations on polyacetylene have primarily employed material made

using this catalyst formulation with a few notable exceptions. Wnek et al. have used a

modification of the Shirakawa conditions to obtain polyacetylene as a foam, whose

electrical properties are similar to those of polyacetylene film, but whose physical form

presents interesting possibilities for potential applications [108]. Wegner's group has

made extensive use of the Luttinger catalyst, a mixture of Co(NO3 )2 and NaBH4 [109].

This latter procedure also produces polyacetylene in the form of films, and as discussed

below Wegner has used results obtained with these materials to challenge the accepted

models for both the stucture and mechanism of charge transport in polyacetylene.

However, the question of whether the Shirakawa and Luttinger polymers exhibit identical

morphologies remains open. Finally the synthetic approach of Edwards and Feast

deserves mention (I 10]; using a polymeric precursor, these authors have attempted to

introduce the conjugated double bonds of polyacetylene via a thermal extrusion reaction.

Although the material produced was similar to polyacetylene spectroscopically, the

decomposition of the starting polymer was apparently incomplete. Nevertheless, this

general approach is probably worthy of continued examination.

Considering both geometric and rotational isomers, there are four possible planar

structures (I-IV) for the polyacetylene backbone. Based on infrared and Raman

assignments and comparisons with model compounds, Shirakawa has found that

. .t - . - r t.. .o.~~~~~.....""............................ "" ..... .....

24

Ziegler-Natta polymerization at -78*C Produces exclusively the cis-transoid structure (1)

[15]. Polymerization at 150'C, on the other hand, yields only the trans-transoid product

(II). Syntheses carried out at intermediate temperatures provide mixtures of I and HI. The

cis-transoid structure may also

7 7r

- --.

[, '- -"

nn

be converted to trans-transoid by heating at 200' C for one hour [111. Although the

trans-cisoid structure (Ill) may be an intermediate in this thermal isomerization, there is

no evidence for its independent existance. Structure IV, the cis-cisoid isomer, is also not

observed (and in fact cannot be rigorously planar for more than two repeat units).

Adopting the common usage we will refer below to structures I and 11 as cis and trans,

S- respectively.

Based on X-ray measurements Shirakawa has estimated polyacetylene produced with

the Ziegler-Natta catalyst to be approximately 85% crystalline but highly disordered

-|-- . .-. ?.-.

[' ""[115. Thiodirerizaionadto broadnio the X-ray ha d iffraon thetaks-tand rodul t

makes c it qiedifiult tay deemnahecytlsrutrsfethrolae smr

6- o

25

Using a combination of the observable diffraction peaks, molecular packing considerations,

and model compound geometries, Baughman has proposed structures for both cis- and

trans-polyacetylene [104,113,114]. Significantly, the chain axes in both structures are

predicted to lie along the fiber axis. This orientation was also suggested by early electron

diffraction measurements [1151. Recently, however, Wegner has interpreted a more

complete analysis of electron diffraction data in terms of a chain-folded structure similar

to that observed in polyethylene and other organic polymers [1161. On the other hand,

Chien and Karasz have also carried out an extensive analysis of electron diffraction data

which they believe supports the Baughman model [1171. Obviously the resolution of this

problem is critically important to the interpretation of the wide variety of physical

measurements which have been carried out on polyacetylene.

An equally important unresolved question involves the extent of crosslinking in

polyacetylene. In an early 13 C NMR experiment Waugh observed approximately 5% sp 3

carbons, which he suggested might represent crosslinking sites [ 118]. However,

subsequent NMR experiments have been unable to detect any sp 3 carbons, placing an

upper limit (essentially the experimental limits of detection) of I% on the extent of

crosslinking (119]. Of course, one crosslinking site for every 100 carbons would still

represent an extensive perturbation of the polyacetytene structure. More recently, Wegner

*0 has suggested that although the (CH) x produced by Ziegler-Natta catalysis is highly

crosslinked, polyacetylene synthesized using the Luttinger catalyst is virtually free of

crosslinks [109]. No in depth investigation of the differences in spectroscopic and

transport properties of these two forms of polyacetylene has yet been carried out.

26

-. This detailed examination of the structure and properties of polyacetylene was, of

course, prompted by the finding that the conductivity of polyacetylene could be raised to

quite high levels by exposure to oxidizing or reducing agents 1120]. Whereas cis- and

trans-polyacetylene have conductivities of -10-9 ohm"t -cm"1 and 10-5 ohm"I-cm"1 ,

respectively, treatment of these films with, for example, AsF 5 leads to conductivities as

high as 1000-2000 ohm' 1 -cm "1 . (For better or worse, in analogy with more traditional

semiconductor terminology this process has generally been referred to as "doping".)

Lightly doped samples exhibit intermediate conductivities. At room temperature the

conductivity of heavily doped samples shows a metallic temperature dependence, but at

lower temperatures the conductivity is activated [121]. Oxidized polyacetylene is a p-type

conductor while reduced polyacetylene is n-type (122,123]. Oxidation of stretch-aligned

films (draw ratio of three) results in a conductivity of approximately 3000 ohm'-cm1

along the direction of orientation while the conductivity perpendicular to the stretching

axis is a factor of 16 lower [106].

After considerable initial debate as to the actual mechanisms of the doping reactions,

it is now generally accepted that all dopants can be divided into three general groups. The

first type consists of oxidizing agents like arsenic pentafluoride, iodine, bromine, and

certain transition metal salts. These dopants function by removing electrons from the

*O polymer r. system to produce a delocalized cation or polycation [124]. In contrast, the

- second group of dopants is composed of reducing agents, such as sodium metal or sodium

naphthalide, which act by donating electrons to the polymer ir orbitals to generate

delocalized anionic species. Oxidation and reduction may also be achieved

electrochemically [125,126]. X-ray diffraction studies on the oxidized polymer indicateS . *

27

that the dopant moieties physically intercalate between the (CH) x chains in heavily doped

samples [127,128].

The third type of polyacetylene doping involves strong proton acids. Gau et al.

initially showed that exposure of (CH), films to the vapors over perchloric or sulfuric acid

led to conductivities as high as 1000 ohm'1 -cm"1 [129]. There was, however, some

question as to whether the acids themselves were the true dopant species, since the

anhydrides of both of these acids are quite strong oxidants and would be expected to be

effective oxidants of polyacetylene in their own right if present in trace amounts in these

experiments. This question became even more significant when it was shown that

hydrogen bromide and hydrogen chloride, acids without oxidizing anhydrides, are at best

weak dopants of (CH)x in the gas phase. The reality of proton acid doping was

eventually demonstrated using trifluoromethanesulfonic acid, whose anhydride can be

shown not to dope (CH)x (130], and anhydrous hydrogen fluoride [131]. Each of these

' ispecies proves capable of doping polyacetylene to metallic levels.

The precise mechanism of proton doping remains undetermined, but logical

speculation suggests that the reaction proceeds by addition of a proton to a double bond.

S- This new sp 3 position then splits the chain into two segments, one a shorter length of

neutral polyacetylene and the other a delocalized cation like that produced in normal

oxidative doping. Using compensation experiments McQuilan has detected the presence

of sp 3 infrared absorptions in HF doped samples [131]. These sp 3 sites would appear to

create a serious disruption of the original extended ir conjugation of polyacetylene,

-'..- particularly at the high doping levels used to obtain metallic conductivities. Just how such

a system could support effective electrical transport remains unclear.

28

One shortcoming of the above doping techniques is that they do not allow specific

reaction of only selected areas of a given polyacetylene film. Because of the highly

porous nature of (CH),, gas phase or solution doping using a conventional masking

technique cannot be used to generate conducting or semiconducting patterns in a film.

Attempts at using ion implantation to dope polyacetylene have provided mixed results.

Recently, Clarke et al. have provided an alternate approach to this problem using

diaryliodonium or triarylsulfonium salts [132]. These materials are inert toward

polyacetylene under normal conditions and may be impregnated into (CH), films without

affecting their conductivity. On ultraviolet irradiation, however, these salts decompose to

yield strong proton acids which, as noted above, are effective dopants of polyacetylene.

In this case simple radiation masks allow selective doping of only those areas of a film

exposed to UV light. This technique also allows one to achieve a desired doping level

quite conveniently by controlling the amount of salt impregnated into the polymer.

Although the chemistry involved in these various doping processes is now fairly well

understood, the means by which these reactions promote high electrical conductivity in

polyacetylene is a source of some controversy. Originally the conductivity enhancement

was interpreted in terms of straightforward removal of electrons from the valence band

(oxidative doping) or addition of electrons to the conduction band (reductive doping) in

analogy to the established mechanisms for the doping of traditional semiconductors. More

recently, however, an alternate explanation invoking a novel soliton theory has been

advanced to explain the doping of polyacetylene. Stimulated by the general current

interest in solitons [1331 in both the theoretical and solid state physics communities, this

theory has received considerable attention and some experimental support.

S.

29

The soliton theory actually attempts to explain two superficially unrelated features of

polyacetylene: the nature of the large number of free spins in undoped trans-(CH)x and

the mechanism by which electrical conductivity is enhanced on oxidation or reduction of

the polymer. Ironically this theory was first postulated in somewhat simpler form in 1962

in an insightful paper by Pople and Walms!ey [134] on possible defect states in extended

polyenes, but was rather quickly rejected by experimentalists. Essentially Pople and

Walmsley suggested that delocalized species like Structure V might contribute significantly

to the high number

V

of free spins observed in conjugated polymers (one spin for every 3000 carbons in

trans-polyacetylene). However, the prediction that these defects should be thermally

populated (Ea~0.2eV) was quickly refuted experimentally [135].

Pople's suggestion was revived in 1979 independently by Rice [136] and by Su,

Schrieffer and Heeger [137], who realized that this defect site could be described as a

soliton domain wall separating two energetically equivalent but topologically distinct

segments of trans-polyacetylene. Although these defects were again predicted to be

accessible thermally, a somewhat higher Ea was calculated (-0.4eV) and the observed

spins were suggested to be defects generated irreversibly during the cis-trans

isomerization. These domain wall defects, often called neutral solitons, are predicted to

be highly mobile along the trans-polyacetylene chain. It is important to understand that

* * * . . . ... - - .-

30

*.. such mobile defects are impossible in cis-polyacetylene, where an equivalent defect would

separate the chain into energetically inequivalent cis-transoid and trans-cisoid structures,

9-." leading to pinning of the defect at a chain end. (In fact few or no free spins are observed

*. in pure cis-polyacetylene [1381, but this is only an indication of a low defect level and not

a consequence of any soliton-like behavior.)

The neutral soliton model predicts that the mobile spins in trans-polyacetylene

should show motionally-narrowed EPR signals with little or no resolved hyperfine

coupling. Experimentally these features are, indeed, observed [139,140]. Moreover,

Nechtschein has used dynamic nuclear polarization (DNP) experiments and proton TI

measurements to show not only that the spins in trans-polyacetylene are moving very

rapidly (_1013 rad/sec), but that their motion is extremely one-dimensional, with a ratio

of the diffusion rates along and transverse to the chains greater than 106 [141]. Although

the absolute magnitude of the spin diffusion rate along the chain has been questioned

[142], the observation of an Overhauser effect in the DNP experiments certainly places a

lower limit of -5x 1010 rad/sec on the diffusion rate [1411. These results provide the

strongest experimental support for the presence of neutral solitons in undoped

, polyacetylene, although Etemad has recently suggested that the soliton theory may also be

necessary to explain phototransport effects in trans-polvacetylene [143,144].

0

The more controversial aspect of the soliton theory involves the actual mechanism of

- polyacetylene doping. According to this theory initial reaction takes place at neutral

soliton sites to provide charged defects (Via for oxidation, VIb for reduction) which also

grn.'

'*~~~~~~~~~~~'-~~~~.7 *-7-*'~ *~~ ~ :<.. -~ *~*i-C 2 ..- 7~

31

I-o

may be described as soliton domain walls [136,137). However, the relatively low number

of neutral solitons allows this mechanism to account only for the first -0.03 % of doping.

The important prediction of the soliton mechanism is that subsequent doping also serves to

generate charged solitons from defect free regions of the polyacetylene chain.

Energetically this is favorable because it only costs 0.4eV to create a soliton (neutral and

charged solitons are of equal energy), while removal of an electron from the valence band

or addition of an electron to the conduction band costs 0.7eV, one-half the band gap

- [137]. These charged solitons initially form states at the band center which are localized

by the coulmbic attraction of the dopant counterion. Ultimately screening occurs and the

solitons form a band at the gap center. This stage would correspond to the

* semiconductor-to-metal transition observed at approximately 1% doping. Continued

doping would cause broadening of the soliton band and eventual overlap with the valence

and conduction bands. Because the charged solitons are spinless, a unique prediction of

the soliton mechanism is that doped polyacetylene should show metallic behavior above

the semiconductor-to-metal transition, but that there should be no Pauli susceptibility

observed until overlap of the soliton band with the valence and/or conduction bands

occurs [145]. At this point the system can be described by a typical single particle band

picture and the usual properties of a metal, including Pauli susceptibility, should be found.

. ... ,• ,; i , € .-... *. . -...

32

The extent of experimental support for the soliton doping mechanism has been the

source of considerable controversy. Ikehata et al. have observed that, under very

carefully controlled conditions, AsF 5 doping leads to a loss of essentially all Curie spins

by 1% doping accompanied by very low levels of Pauli susceptibility until approximately

7% doping, at which point a rather dramatic increase in Pauli susceptibilty is observed

[145]. These results are, of course, quite consistent with the predictions of the soliton

model. On the other hand, in their attempts to repeat the Ikehata experiments,

Tomkiewicz et al. observed a continuously increasing Pauli susceptibility as a function of

dopant concentration, even in samples below the 1% doping level [146]. These results

were interpreted in terms of inhomogeneous doping, where, for one reason or another,

metallic "islands" are formed at very low doping levels and grow in size and/or number as

doping progresses [147]. The semiconductor-to-metal transition is then better understood

as a percolation threshhold. AC conductivity measurements [148] and the non-ohmic

behavior of the samples in electric fields [149] support this picture of inhomogeneity in

*the doping process.

Initially it was suspected that the difference between these two apparently

contradictory sets of results lay in subtle differences in sample preparation. Joint

experiments have revealed, however, that the actual difference involves the measurement

techniques employed [150]. Both groups obtain similar results throughout the doping

range when using X-band EPR spectroscopy to measure the susceptibility. Above 1%

doping there is a systematic deviation between results from the Schumacher-Slichter

technique (used by Ikehata et al. [1451 and those from X-band EPR measurements (used

by Tomkiewicz et al. [146,147]). Since skin depth problems become increasingly severe

in the EPR measurements as the doping progresses above 1 %, one would expect the:0i~

3314

Schumacher-Slichter results to be more reliable in this region. Although a complete

resolution of these discrepancies has not yet evolved, several conclusions can be drawn at

this point. There is certainly a dramatic transition in the Pauli susceptibility above 7%

doping, consistent with the predictions of the soliton mechanism. On the other hand,

there is a finite, measureable Pauli susceptibility observable in samples doped to

concentrations below 7%. In particular, this is true for samples in the 0.2-1.0% range,

below the semiconductor-to-metal transition. This Pauli susceptibility undoubtedly reflects

the nonuniformity of the doping process, even under conditions designed to minimize,

-" inhomogeneity. The extent to which this nonuniformity plays a role in the transport

4properties at low doping levels remains unresolved.

In another experimental probe of the magnetic properties of (CH)x , Peo et al.

examined the magic angle 13 C NMR spectrum of polyacetylene as a function of AsF5

doping (151]. They saw no indication of a Knight shift of the carbon signal by carrier

spins until 7% doping, although conductivity measurements clearly showed a

semiconductor-to-metal transition in their samples at a doping level of roughly 1 %.

- Subsequent experiments by Clarke and Scott, however, have revealed the presence of a

shifted peak in AsF 5 doped samples at doping concentrations as low as 1% [152]. They

attribute the bulk of the shift to a chemical shift arising from the removal of electrons

from the polymer fr-system in the doping process, with any Knight shift representing a

* smaller contribution superimposed on the chemical shift. These latter results also

demonstrated for the first time a difference in the degree of homogeneity of doping

between cis and trans polyacetylene. Although trans samples appear somewhat

inhomogeneous, the cis samples at, e.g., 3% doping appear to consist of heavily doped

.4

Jd

r e. .° .. ..". .

regions and essentially undoped regions. This observation suggests that comparisons

between lightly doped cis and trans samples should be made very cautiously.

Infrared spectroscopy has also provided a measure of support for the soliton doping

mechanism. Rice and Mele have carried out extensive calculations of the infrared features

expected in a polyacetylene sample containing nonoverlapping charged solitons [153].

Their predictions are in quite good agreement with the infrared spectrum of lightly doped

polyacetylene (154]. Interpretation of the infrared data must, however, take into account

the remarkable similarity in the IR spectra of lightly and heavily (10%) doped

*- polyacetylene (154,155], even though the soliton model is not expected to be applicable at

high doping levels.

,5.J

B. Electronic Properties

1. BAND STRUCTURE

Even before the highly conducting properties of doped (CH)x films were discovered,

there was considerable interest in the band structure of this prototype infinite polyene

system [12]. This interest arose from speculation on the physical properties to be

expected from infinite polyene chains. The methods of calculation applied to this system

have addressed the issue of the adequacy of a single particle band description and the

merits of including electron correlation effects. Unfortunately, widely different results

C have been reported for these calculations. With the discovery of the metallic properties of

." doped (CH),, band structure calculations have received renewed attention stimulated by

the desire to obtain a theoretical basis for understanding the electronic properties of both

intrinsic and doped material. A number of band structure calculations appear in the

recent literature on both the trans- and the cis-(CH)x conformations.

5/

5*.** 5, S

35

Calculations of the single particle band structures of cis-transoid and trans-(CH)x

were made by Grant and Batra (156] using both the LCAO extended tight-binding method

and the semiempirical extended Hackel method. For the cis-transoid calculations the

crystal structure proposed by Baughman et al. (114] was used. For trans-(CH), different

carbon-carbon bond length schemes were used including uniform bonds and both weakly

alternating and strongly alternating bond lengths. The cis-transoid isomer was found to

have a band gap of 1.2eV and conduction and valence band widths of nearly 6eV for

directions parallel to the chains. For the trans-(CH)x the degree of bond alternation has a

profound effect on the calculated band gap. For uniform bonds the Fermi level passes

through a degeneracy point at the Brillouin zone edge and the system is intrinsically

metallic. With increasing bond alternation a single particle band gap exists which

increases to 2.3eV for typical single and double lengths.

Band structures for the trans, cis-transoid and trans-cisoid conformations have been

calculated by Kasowski et al. (157] using an extended muffin tin orbital method. This

study shows that (CH)x is an intrinsic semiconductor with energy gaps of 1.6, 1.7 and

1.2eV, respectively, for the trans, cis-transoid and trans-cisoid structures.

* Three-dimensional charge density contour plots show that the bonding orbitals are

C-ir-like, sticking out from the plane of the molecule. These ir-states, which are

responsible for the semiconducting gap, are also expected to react with dopants.

0

Yamabe et al. [158] have compared the calculated band structure properties of the

cis-transoid and the trans-cisoid structures of (CH) x . From total energy/unit cell, ir-bond

order and interatomic interaction energy comparisons these authors conclude that the

cis-transoid isomer is the favored cis-structure of (CH) x.

,. . .. .. . . . . . . .. . . . . . . .-.. o .o .•° . . .

L- 36

The band calculations on heavily doped trans-(CH), by Kasowski et al. [159] are of

considerable interest. Extended muffin-tin orbital calculations were made on defect-free

material and material heavily doped with AsF 5 , AsF , SbF and PFf. For the

hexafluoride dopants, AsF6 and SbF6, hybridization of metal s-states with polymer

ir-states produces a partially filled metallic band implying strong metallic conductivity

along chains and weaker metallic conductivity between chains. Doping with AsF 5 yields a

semiconductor with a completely filled As s-valence band and an empty ir-conduction

band. Doping with PF6 yields an empty P s-band, but charge transfer to F atoms partly

empties a ir-band implying metallic conductivity along, but not between, chains. The

results for AsF 5 and AsF6 doping are particularly important since they give theoretical

support for AsF6 as the effective dopant in highly conducting trans-(CH)x. On the other

hand, many dopants which would not be expected to undergo this sort of hybridization are

also effective dopants of polyacetylene.

2. SEMICONDUCTOR PROPERTIES

Unlike the polymeric metals we have described so far in this chapter, undoped

0 (CH) x is an intrinsic semiconductor. Although it was the metallic properties of the

heavily-doped material which stimulated recent interest, the semiconducting properties of

intrinsic and lightly-doped (CH) x are of considerable interest. In particular the

mechanism of doping and the proposed soliton transport models are areas of great

K::: importance and considerable controversy. Junction studies in which (CH), is the active

semiconducting element also place strong emphasis on the properties of intrinsic and

lightly-doped material.

.-',.... .". . . .. ... . .. .. . o" ".b

r. Wb

37

Band structure calculations described in the previous section indicate that both cis

and trans isomers of (CH) x should be relatively wide band semiconductors with band gaps

in the range of 1.2 to 1.7eV. Photovoltaic threshold measurements by Tani et al. [160]

on In/(CH)x Schottky barrers have yielded a value of 1.49eV for the single particle band

gap of the trans-isomer. Optical absorption [107], photoconductivity [160] and

photoelectrochemical cell studies [161] also indicate a band gap of this magnitude.

Electrical conductivity measurements on (CH) x films doped with halogens and

arsenic pentafluoride, AsFs, were first reported by Chiang et al. [16]. The room

temperature conductivity of the undoped trans-isomer is about 2x I0-5 ohm' t cm"I . For

the cis-isomer, the room temperature conductivity is considerably lower with a value of

about 2x 10-9 ohm-'cm-1 . The conductivity of both isomers increases on exposure to a

wide range of dopants. However, doping of the cis-isomer results in isomerization to the

thermodynamically stable trans form [15,1111 so that lightly-doped initially cis films

become heterogeneous systems. With heavier doping, induced isomerization can be

detected by optical absorption, thermopower and NMR second moment techniques.

Plots of Ina versus I/T for undoped and lightly-doped trans-(CH)x show

activated-like behavior but with some curvature, which probably indicates a distribution of

activation energies [162]. For undoped trans-(CH)x the slope at room temperature yields

an activation energy of 0.3eV. This activation energy is relatively insensitive to dopant

concentration up to about 0.1 mole % and then decreases rapidly as the material

undergoes a semiconductor/metal transition at dopant concentrations near 1 mole %.

4 >o . ." - .. .. . . . .,

L38

Thermoelectric power (TEP) measurements by Park et al. [162,163] and by Kwak

et al. [121,1641 are in good agreement for the entire dopant concentration range from

undoped semiconductor to heavily-doped metallic behavior. For the undoped trans-isomer

a value of about +90 uV/K is obtained. For acceptor dopant concentrations below

0.1 mole % there is little change in thermopower. As the dopant concentration is

increased from 0.1 mole % to 1 mole % the thermopower decreases drastically to a

value characteristic of the metallic state (about 10 uV/K). The relative insensitivity of

TEP to concentration below 10-1 mole % has been cited as evidence for a defect or

. -. impurity concentration of about 0.1 mole % in pristine (CH) x. At low dopant

*- concentrations, Park et al. [162] find that TEP is essentially temperature independent

"" -"which is interpreted in terms of a dilute concentration of carriers hopping in a set of

..-: localized states. From the temperature independence of TEP, implying a constant carrier

density of about 2x 1018 cm "3 for undoped trans-(CH),, and from the activated

conductivity, Park et al. [162] conclude that the mobility is activated in the dilute limit

with a magnitude of about 5 x 10-5 cm 2 /V-sec at room temperature. Since in the metallic

regime the mobility is estimated to be about 60 cm 2/V-sec, an abrupt chang, cf 5 to 6

orders of magnitude in carrier mobility is assumed to occur on passing through the

"O"- semiconductor/metal transition at dopant concentrations near I mole %. The low

mobility of the undoped (CH), is unexpected for a wide band semiconductor. The

inference that transport is by a localized state hopping mechanism is argued to be

qualitatively consistent with the proposed soliton doping mechanism discussed in more

detail above.

Measurements of the electric field, E, dependence of conductivity of lightly AsF 5

doped cis-(CH) have been reported recently by Mortensen et al. [151]. In the high field

[. .

[ " ..°

39

regime (at T=4K) the resistance RH was of the form

RH - RHO exp (Eo/E).

In the low-field regime the temperature dependence of resistance RL fits the relationship

RL - RLO exp (To/T)1 / 2 .

These field and temperature dependences are compared with the expected behavior of

granular metals in a concentration regime in which isolated metallic particles are dispersed

"" in a dielectric continuum.

3. METALLIC PROPERTIES

The metallic properties of heavily-doped (CH) x films have been more extensively

investigated at this time than the semiconducting properties. This emphasis on the highly

conducting state was a consequence of interest in the highly conducting polymers and the

quasi one-dimensional organic charge transfer salts. This interest was further sharpened

when doped (CH) x was shown to have the highest conductivity of any organic material

excluding intercalated graphite. Considering the matted fibrous structure of (CH), films

and the highly reactive nature of the dopants involved, it is remarkable how reproducible

transport results have been for experiments in a number of laboratories. The earliest

indications of high conductivity in heavily doped (CH) x were obtained in the experiments

* of Berets and Smith (14] in which (CH) x powder was exposed to strong oxidants including

the halogens. The pellet conductivity was observed to increase by several orders of

magnitude. However, the discovery by Ito, Shirakawa and Ikeda [15] of a process for

forming continuous films of (CH) x was the key to subsequent research on the electronic

properties of this polymer. In 1977 Chiang et al. [16] reported doping experiments on

[oS

40

(CH) x films showing a metal/insulator transition for an acceptor dopant concentration

near 1 mole % and a maximum conductivity comparable to that of single crystals of the

organic metal TTF-TCNQ. Subsequent experiments demonstrated high conductivity by

exposure to donor dopants (alkali metals) [123] and conductivities as high as

3000 Q2 1cm-1 by using stretch oriented (CH), films [106].

In spite of the magnitude of the conductivity of heavily doped films, which clearly

indicates metallic behavior, the temperature dependence is not characteristic of a metal.

-- Figure 9 shows the temperature dependence of conductivity for highly doped and

intermediate doped trans-(CH), films. In heavily doped material, the conductivity

increases slightly with decreasing temperature for T220K [122]. As the temperature is

decreased further to 4K, the conductivity decreases. However, for temperatures below 4K

the conductivity remains constant [121,164]. For more lightly doped samples the

*conductivity is activated over the entire temperature range.

Thermoelectric power measurements on the highly conducting (CH), films are of

particular interest because they provide a method for evaluating metallic behavior without

2- current flow across barriers associated with the disordered fibrous structure. Kwak et al.

measured TEP as a function of temperature in AsF 5 doped films [121,164]. For greater

than 1 mole % the thermopower was found to be small (z 10&V at 300K), positive and

independent of concentration. The temperature dependence was linear and extrapolated

r." through the origin as expected for an ideal simple metal. The sign of the thermopower is

consistent with the acceptor character of the dopant, and the magnitude and temperature

L dependence is indicative of intrinsic metallic behavior.

41

The Hall effect in heavily doped (CH) x films has been measured by Seeger et al.

[122]. The Hail voltage was found to be very small necessitating double phase-sensitive

techniques to obtain an adequate signal-to-noise ratio. Measurements were made over the

temperature range from 4K to 300K and for acceptor concentrations (AsF 5) from 2

mole % to 17 mole %. The sign of the Hall coefficient RH is positive and relatively

insensitive to temperature (RH increases by a factor of 2 or 3 on cooling to 4K). The

magnitude of RH is very small, which, if interpreted by a simple single band model, yields

a carrier concentration two orders of magnitude greater than the dopant concentration.

Since in the single band case RH=I/ne(PH/p) where T H is the Hall mobility and A is the

conductivity mobility, the data imply that AH; 10- 2;L if the carrier density is assumed

equal to the AsF 5 dopant density.

4. JUNCTION PROPERTIES

The intrinsic semiconductor nature of (CH) x together with the ability to dope the

material with either donor or acceptor molecules forming n- or p-type material has

generated interest in junction properties and their possible device applications. The first

reports on junction properties of (CH) x described rectification properties observed at a

crude junction formed by mechanically pressing together sheets of n- and p-type (CH) x

[81]. Subsequently a number of studies have been reported involving Schottky barriers

using heavily doped (CH), as the metal in contact with inorganic semiconductors [165],

Schottky barriers in which the (CH) x is the semiconductor [160,1651, heterojunctions with

inorganic semiconductors [165] and electrochemical cells [161] using (CH) x as the active

photoelectrode.

7 -77-; ~i .- 7 -'- -7

42

Schottky barriers were formed by polymerizing (CH), directly on n-type

semiconductors (Si and GaAs) and subsequently heavily doping the (CH) x with AsF 5 to

make it metallic (165]. These junctions had barrier heights in the range from 0.7 to

1.0 volts as estimated from I-V and C-V characteristics. Junction formation showing the

p-type conductivity of undoped trans-(CH), was achieved with pressed contacts of low

electronegativity metals (Na, Ba, In) [165].

Extensive studies of Schottky barrier characteristics with undoped and lightly-doped

(CH), as the active semiconductor have been reported by Tani et al. [160]. Initial

experiments were directed toward reliable determination of the band gap of rans-(CH),

by measurement of the photovoltaic effect threshold. With In (q)M= 4 .12eV) as the low

work function junction metal and Au (M-5. leV) as the ohmic contact, cells with

excellent Schottky barrier I-V and C-V characteristics were obtained. -From the

photovolatic response threshold a value of 1.48eV was obtained for the single particle

band gap in trans-(CH)x . The band gap was also shown to be independent of doping

level for a conductivity range from 2.6x 10-4 to 4 2'cm-1 .

The junction characteristics of a series of In/(CH), Schottky barriers with a range

of doping concentrations have been used to determine mobility as a function of

conductivity for AsF 5 -doped trans-(CH)x [166]. From the linear slopes of 1/C 2 versus

V, characteristic values of the ionized acceptor density NA were obtained for each (CH)1

conductivity. Assuming that the carrier density equals NA, mobility values ranging from

03 x 10-6 cm 2 /V-sec for undoped (CH)x to 3 x 10-4 cm2 /V-sec for lightly-doped.3

(a- I x 10-4 Q-Icm-1 ) material were obtained. These values are consistent with mobilities

,. .* inferred from doping concentrations by Park et al. [162].

.~~~~~~~~~ 7- . - .** ** * * * * -. ?L33 - * . . . .e *

43

Heterojunction formation using p-type undoped trans-(CH), in contact with n-ZnS

has been demonstrated by Ozaki et al. [165]. An open circuit photovoltage of 0.8 volts

was observed with significant photosensitivity in the energy range from 1 to 3eV

attributed to carrier photogeneration in the (CH) x. Another potential solar cell device in

which (CH), is the active photoelectrode in an electrochemical cell has been described by

Chen et al. [161]. This cell consists of undoped (CH), and Pt as electrodes in a sodium

polysulfide electrolyte. An open circuit photovoltage of 0.3 volts was obtained with

approximately one sun illumination on the (CH) ,; however, cell efficiency was limited by

series resistance and the small effective area of the (CH), electrode.

V. OTHER CONDUCTING POLYMERS

Initially it was hoped that substitution of the hydrogens of polyacetylene with

various functional groups would allow the constuction of an entire family of conducting

polymers whose properties could be controlled in a systematic way. As in the case of

(SN), these expectations have not been completely fulfilled, although some progress has

been made. Straightforward derivatives of (CH), where alternate hydrogens have been

replaced by methyl [167] or phenyl [168] groups, for example, can be synthesized as black

powders, but can not be doped to high levels of conductivity with either oxidizing or

reducing agents. The failure of these substituted materials to dope to high levels of

conductivity is not yet understood. Space filling molecular models suggest that even in the

less crowded trans form of polyacetylene, substituents which are sterically more

demanding than hydrogen (or perhaps fluorine) encounter serious van der Waals

repulsions if the polymer backbone is forced to remain planar. Twisting of the backbone

relieves the steric interaction, but also diminishes the overlap of the ir orbitals along the

chain, an effect which would certainly change the electronic properties of the polymer.

S .*% bb.*bI

44

On the other hand, these polyacetylene derivatives probably also differ from (CH) x itself

in terms of molecular weight, morphology, crystallinity, and other parameters whose

effects on the doping process are not currently known.

One attractive approach to lessening the steric interaction in polyacetylene

derivatives involves the copolymerization of acetylene itself with a substituted acetylene.

Chien and co-workers [169] have shown that acetylene and methylacetylene can be

copolymerized with the proper choice of catalyst to produce films which can be doped

with AsF 5 to conductivities as high as 50 ohm'l-cm "I . Varying the ratio of acetylene to

methylacetylene in the polymer produces interesting changes in both the electronic and

mechanical properties of these films. Thus the ultimate conductivity achieved on AsF 5

doping seems to decrease as the fraction of methylacetylene in the copolymer increases.

On the other hand, for compositions containing greater than 50% methylacetylene films

are produced which show elastic properties under certain conditions. When dry these

films are relatively brittle, but on wetting with solvent they can be stretched quite easily

and retain this elongation if dried under tension. Rewetting of one of these stretched

films then causes it to revert to its original shape. Unfortunately the copolymers

synthesized to date appear to be even more sensitive to oxygen than (CH),.

A more highly substituted derivative of polyacetylene has also been reported.

Gibson et al. [70] have polymerized 1,6-heptadiyne under homogeneous Ziegler-Natta

-: conditions to produce free standing films with a green-gold metallic luster. Spectroscopic

measurements indicate that poly( 1,6-heptadiyne) has a conjugated backbone with

6-membered rings fused 1,3 along the chain (VIIa).0,

% .......... ,..-*.*.*.% \..................**. ** .

*- ---. ... -, .

45

Unlike polyacetylene this material is normally amorphous and continuous rather than

fibrillar. On doping with iodine the conductivity rises from 10-12 ohm-1 -cm-1 for the

pristine polymer to approximately 10- 1 ohm- I-cm- for a material of composition

,.(C 7 H8 11 0 ); treatment with AsF 5 leads to a maximum conductivity of 10-2 oh4-.5cm'.

* [(1711. The ultimate limitation on the conductivities achieved on doping appears to lie not

in the intrinsic electronic properties of the polymer, but rather in the facile rearrangement

* of the conjugated backbone structure VIla to the cyclohexadiene structure VIM, a process

which interrupts the conjugation along the polymer chain. This rearrangement is

apparently catalyzed by the same species which dope the polymer. Thus the reaction with,

for example, AsF5 actually consists of two competing reactions, a faster oxidation which

* raises the conductivity of the polymer and a slower rearrangement which reduces the

conductivity. This problem eventually limits the applicability of poly( 1 ,6-heptadiyne), but

the simple demonstration that this (CH), derivative can be doped to reasonably high

* conductivities should encourage the search for other conducting derivatives.

MacDiarmid and his co-workers have taken a somewhat different approach to the

0 synthesis of (CH), derivatives [172]. Rather than polymerizing substituted acetylenes,

* . they have produced new materials by carrying out reactions on polyacetylene itself. In

particular they have shown that on heating, bromine-doped (CH-)x evolves H~r and Br2 to

give golden semiconducting films of composition (CHqBr y)x where (q+y)m0.82-0.98.

Thus a sample of composition (CH 0 .8 2 Br0.14 )x showed a conductivity of

S "1

..............................

-. . . .. . . . . . . . . . . . . . .

- 46

2.1 x 10-6 ohm'l-cm "I . As with (CH) x itself, subsequent oxidation of these brominated

films with AsF 5 or iodine then leads to dramatic enhancements in conductivity with

ultimate values on the order of 10 ohm'l-cm 1 being achieved. The actual structure of

the brominated polymer has not been determined, but in addition to formal replacement of

hydrogen by bromine some acetylene formation and crosslinking are believed to occur.

One additional polyacetylene derivative has received considerable theoretical

attention, although its actual synthesis has not yet been achieved. Band structure

calculations suggest that poly(difluoroacetylene), (CF) x , should be quite similar to (CH) x

electronically, with the exception that all of the (CF)x bands should be significantly lower

in energy than those of (CH) x [173]. This suggests that (CF) x might provide a more

stable n-type material than does (CH) x on treatment with reducing agents. Unfortunately

the difluoroacetylene monomer is highly unstable and more elaborate synthetic routes will

be required for the generation of this interesting polymer.

Rather than pursuing direct modifications of polyacetylene, several research groups

have begun examinations of other polymer systems with conjugated backbones. The first

successful demonstration of high conductivity in a nonpolyolefinic organic material

involved poly(p-phenylene) (PPP).

~j. PPPPrevious workers had shown that PPP forms charge-transfer complexes with oxidizing

agents such as iodine (1741, but the conductivities of these complexes never exceeded

°. . . . '~. . '

47

-4x 10-5 ohm' 1 -cm " 1. Using AsF 5 , a considerably stronger oxidizant, Ivory et al. [175]

were able to achieve conductivities as high as 500 ohm'l-cm- l . In comparison the

conductivity of undoped PPP is less than 10-14 o _m-cm1. It was also shown that

reducing agents such as sodium naphthalide or various alkali metal vapors could be used to

generate conducting n-type PPP. The final conductivity of the doped polymer can be

varied by controlling the extent and duration of exposure of the PPP to the dopant

species. As in the case of polyacetylene, more lightly doped samples behave as

semiconductors. Junctions made by pressing together pellets of semiconducting p- and

n-type PPP have been shown to exhibit the rectifying characteristics expected for a p-n

junction. Hall measurements on heavily AsF 5 doped samples provide an estimate of

-1 cm2/volt-sec for the carrier mobility in these systems. (It is interesting that in PPP

reasonable mobilities can be obtained from a straightforward Hall analysis assuming one

carrier per dopant molecule, whereas in polyacetylene similar analyses yield physically

unrealistic mobilities, presumably because of the complications introduced by the fibrous

nature of (CH), films [122].)

In terms of fabricability, poly(p-phenylene) represents less an advance than an

alternative to the possibilities offered by polyacetylene, which is essentially limited to films

grown on the surface of a Ziegler-Natta catalyst solution as discussed above. PPP as

synthesized is an insoluble powder which does not show a glass transition temperature or0

melting point; however, it is susceptible to high temperature sintering, a powder

metallurgical technique which allows the molding of undoped PPP into a variety of shapes.

Moreover, Baughman and co-workers have recently developed a unique synthesis of AsF 5

doped PPP which which may provide an alternative to this sintering process [176].

Oligomers of poly(p-phenylene), for example biphenyl, terphenyl, and quaterphenyl, form

%°

48

complexes with AsF 5 which on prolonged exposure to higher pressures of AsF 5 (-400

tor') are converted to highly conducting materials. Various data suggest that the reaction

involved is actually a simultaneous oxidation and- polymerization which generates a

.."product quite similar to that obtained by AsF 5 treatment of PPP. More importantly, when

single crystal plates of terphenyl are reacted with AsF 5 the resulting polymer exhibits

* strong optical and electrical anisotropy; this type of ordering has not been possible to

achieve using PPP itself. Thus the promise of these solid state polymerizations is quite

exciting.

Obviously if polyacetylene and poly(p-phenylene) can each be doped to high levels

of conductivity, the copolymer consisting of alternating phenylene and double bond units

should be an interesting material. Wnek et al. have, in fact, examined the properties of

this material, poly(p-phenylene vinylene) (PPV) [177].

40CH:=CH }

Pressed pellets of PPV show a conductivity of -10- 10 ohm'l-cm " 1 at room temperature.o

Exposure to 30 torr of AsF 5 for 2-3 hours leads to a material of composition

"C8 H6 (AsF 5 )0 .9 2 ]x which shows a conductivity of -3 ohml-cm tl . The materials

* examined to date have all been low molecular weight oligomers (degree of polymerization

7-9), primarily due to the insolubility of PPV, which causes its precipitation at early stages

/° of the synthetic reaction. The conductivities observed are quite high for such low

0molecular weights, however, and may indicate that further polymerization is occurring

during the AsF 5 exposure as in the case of quaterphenyl.

'

"70

49

Perhaps the most interesting of the polyphenylene derivatives investigated to date

4 are the heteroatom linked phenylene polymers shown in structure VIII.

.--.

In the case where X is oxygen or sulfur these materials are known to exhibit good melt

and solution processing characteristics, features which are seriously lacking in the

polymers described above. This combination of processibility and the possibility for

extended conjugation led two different research groups to independently investigate the

doping of these materials [178,179]. Poly(p-phenylene oxide) (X=O) reacts extensively

- o..." with AsF 5 but does not produce a material with high conductivity. Treatment of

poly(p-phenylene sulfide) (X=S) with AsF 5 , however, leads to conductivities of

1-10 ohm 1 -cm1 in the resulting polymer.

Poly(p-phenylene sulfide) is a commercially available material [180], which exhibits

excellent thermal and chemical stability in addition to the processibiity which allows its

* generation in the form of film, fiber, and monofilament. One must take advantage of

these properties before doping, however, since the reaction with AsF 5 also causes the loss

of many of the desirable properties of PPS. The polymer becomes increasingly brittle asO

the level of doping is raised. The ready processibility of the pristine PPS disappears on

.'-" doping, as the resulting material shows no melting point or evidence of solubility. The

conductivity of the doped polymer is sensitive to heat and exposure to the atmosphere;

either treatment leads to an irreversible decrease in the sample conductivity. Despite

.4?..

- 4 -* ,"* * ° ...

50 ~

these liabilities PPS represents a significant step toward more tractable conducting

polymers.

The doping of PPS has also raised several interesting scientific questions. Perhaps

the most important arises from an examination of the crystal structure of PPS as

determined by Tabor et al. (181]. Succesive phenylene groups along a polymer chain are

twisted at angles of ±450, a situation which at first seems unlikely to lead to the extended

conjugation necessary for conductivity. In their efforts to understand this situation, both

the groups at IBM (1781 and at Allied [1791 have come to the conclusion that on doping

SjPPS is undergoing extensive chemistry in addition to the simple electron transfer proposed

for polyacetylene doping [115,124]. The evidence suggests that coupling of phenyl rings

is occurring to give either cyclized or crosslinked structures, depending on whether the

reacting phenyl groups are on the same or different chains. Indeed, all of the

phenylene-type polymers seem to show such extensive chemical modification on treatment

with AsF 5 that the resulting materials remain poorly understood. In this light the

mechanisms proposed for the reaction of (CH) x with AsF 5 seem quite remarkable in their

relative simplicity.

Heterocyclic aromatic polymers of the general structure IX should also be interesting

* materials from an electronic point of view. In particular, there exists an extensive

literature describing "nyrrole black", a material (or possibly several materials given the

general lack of structural characterization) formed by the oxidation of pyrrole under a

-- variety of different reaction conditions. These powders showed generally high

conductivities and were often assumed to have structure IX where X-N.

* . . . . . . 4 44

" 51

I-I

More recently, Diaz et al. [182] have developed a controlled electrochemical

* synthesis of this polymer, more correctly named polypyrrole, which leads to flexible films

with conductivities as high as 40-100 ohm-l-cm "1 [183]. On the basis of electrochemical

and spectroscopic data the pyrrole polymerization is believed to involve oxidative coupling

- at the a-positions to generate structure IX; simultaneously, however, the polymer is

oxidized electrochemically to provide a material of typical composition

(C4 H3 .4 N0 .9 )4 (BF 4 ). The BF- is incorporated from the supporting electrolyte as a

counterion to the positively charged polymer chain. The polypyrrole is formed on the

electrode surface as a smooth continuous film whose thickness can be controlled by

monitoring the amount of current passed in the electrochemical cell. Thicker films can be

readily peeled from the electrode surface intact. In the absence of air these films can then

be repeatedly cycled between the oxidized and neutral (nonconducting) forms by cycling

the cell potential. Initial exposure to air leads to a reaction of undetermined nature which

causes the polypyrrole to no longer be reducible electrochemically [184]. After the initial

reaction with the atmosphere the resulting polypyrrole films still exhibit the same high

levels of conductivity. In addition these films are remarkably stable, showing no loss in

* conductivity after storage in air for several months. Polypyrrole is also considerably more

thermally stable than the other doped polymers described above. These unique

characteristics should make polypyrrole suitable for a variety of applications without the

elaborate protective measures required for other conducting polymers.

\ *. * . ,. i ., .7. 7 7.. . -....... . ........... .. ........

52

A number of N-substituted pyrrole derivatives have also been polymerized

electrochemically, but their room temperature conductivities tend to be on the order of

10- 3 ohm-l-cm" 1 [185]. Copolymers can also be synthesized readily and show interesting

properties. The conductivity of the copolymer of pyrrole and N-methylpyrrole, for

example, can be selectively changed from 10-3 to 102 ohm-l-cm-1 by varying the mole

fraction of N-methylpyrrole in the electrochemical cell from 0 to 1 [184].

Polythiophene (IX where X=S) has been synthesized by both chemical [186,187]

and electrochemical [188] techniques. As in the case of polypyrrole, the electrochemically

generated polymer is simultaneously oxidized under the reaction conditions. The

chemically prepared polymer is subject to subsequent oxidation by iodine. In neither case

are high conductivities obtained, however. Typical values are in the range of 10- to

10-4 ohm-l-cm " 1. Unlike polypyrrole, the oxidized polythiophene is also unstable in air.

A rather different class of conducting polymers lies outside the scope of this review

but deserves mention. Marks and co-workers [189] and Kuznesof et al. [190] have used

various polymer backbones to hold phthalocyanines in a face-to-face configuration.

* Partial oxidation of these materials with iodine then provides materials with conductivities

. as high as 0.1-1 ohm'1-cm " 1. However, the conductivity in these polymers is not along

the polymer chain itself as in the polymers described above, but is instead similar to that

* observed in other monomeric stacked phthalocyanine charge-transfer salts.

0

.0

C- * .* *. '. - . A

-53

VI. CONCLUSION

A. Potential of Polymeric Metals

Polymers can be made to exhibit semiconducting, metallic, or even superconducting

properties not traditionally associated with these materials. This has encouraged the belief

that conducting polymers may eventually challenge the more classical materials in certain

areas.

It has been suggested that conducting polymers might compete with metals in the

area of light-weight wiring. Of the presently known systems only AsF5 treated graphite

*- appears to have sufficiently high conductivity [191]. Barrier height measurements of

(SN)x films and various semiconductor substrates have shown that (SN) x has a higher

work function than elemental metals [81]. Cohen and Harris [82] have taken advantage

of this to enhance the open circuit voltage of GaAs solar cells. Polyacetylene has also

been shown to form Schottky barriers [160,192] and heterojunctions [51]. Tskukamoto

et al. [193] have fabricated Schottky barrier type solar cells using HCI doped (CH), as

the semiconductive element. The energy efficiency is about 0.2%. Pristine (CH)x has

been used as the photoelectrode in a photochemical photovoltaic cell [161].

In addition to their use in the role of classical semiconductors and metals,

challenging established materials in established technologies, conducting polymers offer

additional technological opportunities which take advantage of their unique properties.

The electroactive properties of the polypyrroles and (CH) x have been explored, and it has

shown that these materials can be repeatedly electrically switched between the metallic

and the insulating states - a change which can be followed optically as well as electrically

[183]. The large surface area of polyacetylene (60 m2 g"1) (194] and the fact that the

54

fibrils can be covered with finely divided silver metal (55] suggest possible catalytic

applications. It is also possible that replacing carbon loaded "conducting" polymers with

intrinsically conducting polymers would lead to superior performance, avoiding the

problems of high frequency cut-off and high field breakdown associated with the carbon

particles.

The metallic properties of conducting polymers permit them to be utilized as

electrode materials. Nowak er al. [195] have shown that (SN) x has good electrode

characteristics, and Diaz et al. [183] have shown that polypyrrole tetrafluoroborate forms

0 a stable nonporous electrode. Noufi et al. [57] have used polypyrrole as a passivating

layer to prevent dissolution of n-GaAs photoanodes in a photochemical solar cell.

Exciting as the above outlined prospects may be, it is important to recognize the

problems that must be solved before conducting polymers can find widespread

technological application.

From the discussion of the magnitude of the conductivities of presently known

~i conducting polymers (except perhaps AsFS-treated graphite), it is obvious that these

materials will not displace copper or the classical metals. In most applications, if they are

to become practical materials rather than laboratory curiosities, their usefulness will most

probably be found in a rather unique blend of their electronic, mechanical and processing

characteristics. Indeed, conducting polymers are probably best considered not as

competitors for the classical metals or semimetals, but as materials providing new

opportunities due to the incorporation into conducting materials of such attractive polymer

properties as melt and solution processibilities, low density and plasticity or elasticity.

V -. "°% -kc~

55

Ideally, these polymers would also satisfy a variety of increasingly important ecological

considerations such as low toxicity and nonenergy-intensive synthesis and processing.

. Preparation from nonpolluting aqueous media would, of course, be desirable. Advantage

- -could also conceivably follow from the thermal conductivities and high absorption

- coefficients of these metallic polymers.

-, . •

Though some of these properties, such as low density and high absorption

coefficient, characterize all of the existing conducting polymers, some of the other

desirable properties described here are not typical of either (SN), or the chemically (as

opposed to electrochemically) modified systems. In particular, neither (SN) x nor the

chemically modified conducting polymers are stable in air (particularly the alkali metal

reduced n-type polymers), nor are they thermally stable much above room temperature.

The majority of dopants used to impart conductivity to the polymers (e.g., AsF 5 , 12, Br 2 ,

etc.) are highly toxic. In addition, the chemical doping process universally appears to

degrade the mechanical properties of these polymers, making them brittle where previously

they were flexible. In contrast to the chemically oxidized systems, electrochemically

oxidized polypyrrole-BF 4 , after initial air oxidation, is stable in air for several months,

though its electrical conductivity slowly decreases, and shows appreciable thermal stability,

though its flexibility is not yet ideal. Some of the copolymer films of acetylene and

-0 methylacetylene reported by Karasz and coworkers have exhibited elastic properties when

wet with solvent (23]. When dry these films are relatively brittle but on wetting again

with solvent they can be readily stretched and retain their elongation if dried under

,tension.

iI_

t~~~~~ + ++

-. t aL,4 . . . .. .-. =.4**...-

-°. * S . * ** CS S .•

56

Obviously there will be certain applications such as batteries where the conducting

polymer system can be maintained in a restricted atmosphere and considerations of air

stability and flexibility are not crucial. However, for widespread applications, taking full

advantage of these materials, a significant amount of polymer engineering will be

necessary to stabilize these materials and improve their polymeric properties in the

metallic and semiconducting state. The increasingly wide variety and complexity of

polymers which can be made to show high conductivity bodes well for the ability of

scientists to incorporate these desirable properties into such materials.

ACKNOWLEDGMENTS

We would like to thank the Office of Naval Research for partial support of this

* . work.

..1-

..

b .

..0' SQ

l " S

S

S"

S- .. ... . . ..- . . . '.'. . ., . . . .... 'S; ; '% , ' . '

"-" -" " '. '-.q "

"%' " """-% """'" "" '""-"%""-""-' '" " " % % '" "'""" ** ,' -' -" - "

-. .w .- .

57

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.4p."

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'~~~~~~~~~.'..;. ............. ........... .'.................-....,- ... . , . . - . -, .- - -..- , -- ,,-,- - , - . - , - . . .

67

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-ZS.

68

150. T. C. Clarke, M. A. Druy, J. Flood, A. J. Heeger, S. Ikehata, A. G. MacDiarmid,

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69

164. J. F. Kwak, W. D. Gill, R. L. Greene, K. Seeger, T. C. Clarke and G. B. Street,

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72

C-iu

(N CL

0 LO CV) CV

0 V-

70

e n o n 0 n LI in04 ( N C11

- IA

/ __________________

/Vl kv 14

040

iii Ci,

L6-

-~ ('K

/j 7 CI

. ~ • . . - ° . . ° - . . - .-. . . - ° -. a ..- .

.aO 974

(a)1.0 RI

(b) .. 3MIN

rTc =O.26*K PR-I 4Ta.~.

4 0

~0.1 j/p T

i i(e)Ii 0~=.31"K --0K 80

* . 0 J i i I nli i ii i I I I I i I I

0.-. RT.

-T 0- K P,40j<o/25" "p T-0.26K0---

(e)~6 MFT 8 .o/

"-"0.0 01 PRT I/ol0." 0. .0 5.0 10.0 50 10 0

T(OK)

Figure 3. Temperature dependence of the resistivity of (SN)x crystals: (a) (SN)x pellets;

(b) crystals from Walatka [5]; (c) crystals from Street [23]; (d),(e),(f) crystals grown by

* technique described by Street and Greene [57].

'9 •°

--:-:.:,,:. .

~ -aa~a.a aa a... a. I. . . . * <* a a % aai a aa . aa a a0.0 0 1 ' a . ,* , 1aa aa h a a a 11 1 . I a I

75

THERMOELECTRIC POWER (MV K-1)

-~2-

m 8rn/

/l 4m

CODCTVT (12 /-1

poe n odctvt f(N,]Fiur 4. Teprtr* eednc ftemeeti

fim gon tvriu sbtat emeatrs T 356/

76

0

Figure 5. Electron diffraction pattern from (a) (SN)x fibers showing the b c reciprocal

lattice set; (b) (SNBrO. 4 )1 fibers oriented similarly to (a); (c) electron micrograph of

(SNBro.4)7 twinned fibers with 20A dimension [9].

77

0.45

0.40 - Pristine (SN) x

o Brominated (SN) xo- Drude-Lorentz Fit

0.35 -%

0.30 -0

R11b 0

0.25

0.20 -

0.15 -

0.05 -

0.14,

1.0 1.5 2.0 2.5 3.0 3.5

0.12

0.10

Rib0.08 -

0.06

0.04 -

0.0210.5 1.0 1.5 2.0 2.5 3.0 3.5 4.0

eV

Figure 6. Polarized reflectance of the sae (SN) 1 crystal before and after bromination

F9.

o W .

78

co

CO

z

C*C-,

U'-U

aouissa pai-wO

79

4000

500 k \

* 400 /

IdI.

Figur30 8.Pesr eedneo tesprodcigtasiintmeauei'S~

[60. he(S~x at i o mM ertal 6.

--,. . . *.. * *.,

o .*°

80

00

01.0

0.60

*. ".

00 °

.- t//,

::0.4

~0.2

00

:-:: /

0 100 200 300T (K)

IFigure 9. Temperature dependence of conductivity for intermediate and heavily doped

[. trans (GCl)x [122,1641.

Table 1. Comparison of shortest interchain bonds in the Mikulski et a,. 21

structure of phase (SN) x with the corresponding Van der Waals Radii.

Interchain Distances In (SN) x

Shortest Within Between V.D.W.Separation (102) (102) (A)

Plane Planes

S-S 3.48 3.72 3.70

N-N 3.35 3.37 3.15S-N 3.26 3.40 3.35

%"U

I,

Ib

I•

I,],"""",":- . . ='""","". ". .- . , .""."',"". .. . ,; . . "." -".-',, "." -. '" ' ,. -' , , , . . '.' L 3' t '', , .. , .. " ' "" "" ' """" "" , ."'' . . . . . . . ,

i Ik

I.,

Ai

~ Yi;l

Il;*

~~Y .~